1. Field of Endeavor
The invention concerns the field of materials science. It relates to a process for producing a single-crystal component or directionally solidified component which is made of a nickel-based superalloy and has relatively large dimensions. Particularly good properties, in particular a very good fatigue strength with low cyclic loading of the component, can be achieved.
2. Brief Description of the Related Art
At high loading temperatures, single-crystal components made of nickel-based superalloys have, inter alia, very good material strength but also good corrosion and oxidation resistance, as well as a good creep strength. On account of these properties, when using such materials, for example in gas turbines, the intake temperature of the gas turbines can be raised so that the efficiency of the gas turbine system increases.
In simplified terms, there are two types of single-crystal nickel-based superalloys.
The first type, to which the present invention relates, can be fully solution annealed so that the entire γ′ phase lies in solution. This is the case, for example, for the known alloy CMSX4 with the following chemical composition (in % by weight): 5.6 Al, 9.0 Co, 6.5 Cr, 0.1 Hf, 0.6 Mo, 3 Re, 6.5 Ta, 1.0 Ti, 6.0 W, remainder Ni; or the alloy PWA 1484 with the following chemical composition (in % by weight): 5 Cr, 10 Co, 6 W, 2 Mo, 3 Re, 8.7 Ta, 5.6 Al, 0.1 Hf; and the known alloy MC2 which, in contrast to the previously mentioned alloys, is not alloyed with rhenium and has the following chemical composition (in % by weight): 5 Co, 8 Cr, 2 Mo, 8 W, 5 Al, 1.5 Ti, 6 Ta, remainder Ni.
A typical standard heat treatment for CMSX4 is, for example, as follows: solution annealing at 1320° C./2 h/shielding gas, rapid cooling with a blower.
The second type of single-crystal nickel-based superalloys is not fully heat treatable, i.e., in this case only a specific part rather than the entire proportion of the γ′ phase enters solution during solution annealing. This is the case for example for the known superalloy CMSX186 with the following chemical composition (in % by weight): 0.07 C, 6 Cr, 9 Co, 0.5 Mo, 8 W, 3 Ta, 3 Re, 5.7 Al, 0.7 Ti, 1.4 Hf, 0.015 B, 0.005 Zr, remainder Ni; and the alloy CMSX486 with the following chemical composition (in % by weight): 0.07 C, 0.015 B, 5.7 Al, 9.3 Co, 5 Cr, 1.2 Hf, 0.7 Mo, 3 Re, 4.5 Ta, 0.7 Ti, 8.6 W, 0.005 Zr, remainder Ni.
The nickel-based superalloys of the second type are usually exposed to a two-stage heat treatment (aging process at lower temperatures) since at higher temperatures, such as are typically used for solution annealing the alloys of the first type, the melting point initiation temperature is already reached and the alloy therefore undesirably begins to melt.
A typical two-stage heat treatment of the alloy CMSX186 is for example as follows:
1st stage: 1080° C./4 h/blower
2nd stage: 870° C./20 h/blower.
The creep strength of the first type of nickel-based superalloys is normally higher than that of the second type, assuming that the alloys belong to the same generation. This is primarily due to the fact that the dissolved γ′ is the main source of the achievable strength.
Nickel-based superalloys for single-crystal components, as are known for example from U.S. Pat. No. 4,643,782, EP 0 208 645, U.S. Pat. No. 5,270,123, and U.S. Pat. No. 7,115,175 B2, contain alloying elements which strengthen the solid solution, for example Re, W, Mo, Co, Cr, and elements which form γ′ phases, for example Al, Ta and Ti. The level of high-melting alloying elements (W, Mo, Re) in the basic matrix (austenitic γ phase) increases continuously as the loading temperature of the alloy increases. For example, standard nickel-based superalloys for single crystals contain 6-8% W, up to 6% Re and up to 2% Mo (in % by weight). Furthermore, small proportions of C, B, Hf and Zr are often present. The alloys disclosed in the abovementioned documents have a high creep strength, relatively good LCF (low cycle fatigue) and HCF (high cycle fatigue) properties and a high resistance to oxidation.
These known alloys were developed for aircraft turbines and were therefore optimized for short-term and medium-term use, i.e., the load time was designed for up to 20 000 hours. By contrast, industrial gas turbine components have to be designed for a load time of up to 75 000 hours or even more.
By way of example, after a load time of 300 hours, the alloy CMSX-4, which is known from U.S. Pat. No. 4,643,782, when it was tested for use in a gas turbine at a temperature of over 1000° C., underwent considerable coarsening of the γ′ phase, which disadvantageously leads to an increase in the creep rate of the alloy.
It is known prior art to subject superalloys of this type to heat treatment after the casting process, in which heat treatment, in a first solution annealing step, the γ′ phase, which is precipitated non-uniformly during the casting process, is completely or partially dissolved in the microstructure. In a second heat-treatment step, this phase is precipitated in a controlled manner again. In order to obtain optimal properties, this precipitation heat treatment is carried out in such a way that the finest possible uniformly distributed particles of the γ′ phase are produced in the γ phase (=matrix).
However, it has been found that directional coarsening of the γ′ particles, the phenomenon known as rafting, disadvantageously occurs in the microstructure of alloys of this type if a mechanical load is present with long-term high-temperature loading (creep loading), or after plastic deformation of the material at room temperature, which is followed by heat treatment (high-temperature annealing) of the material. At high γ′ contents (i.e., at a γ′ content of at least 50% by volume), the microstructure is thereby inverted, i.e., γ′ becomes the continuous phase in which what was previously the γ matrix is embedded.
Since the intermetallic γ′ phase tends toward environmental embrittlement, under certain loading conditions this subsequently leads to a massive drop in the mechanical properties—in particular the yield strength—at room temperature (25° C.) compared to samples which were not subjected to such prior creep loading. This impairment of the yield strength is described by the term “degradation” of the properties (see Pessah-Simonetti, P. Caron and T. Khan: Effect of long-term prior aging on tensile behaviour of high-performance single crystal superalloy, Journal de Physique IV, Colloque C7, Volume 3, November 1993).
A similar effect which leads to the rafting of the γ′ phase also arises during the solidification of nickel-based superalloys on account of dendritic segregations. Particularly in superalloys with a high proportion of elements which diffuse slowly, e.g. rhenium, the segregations of these elements cannot be eliminated fully within an acceptable homogenization time. Since the γ′ phase which precipitates during the cooling has a smaller lattice constant than the γ matrix, but the γ/γ′ lattice offset in the dendrites is greater than in the interdendritic areas, internal stresses are formed during the heat treatment, in particular during the cooling. This results in a change in the γ′ microstructure, in that the initially cubic form of γ′ changes into an elongate form of γ′. This is accompanied by the impairment of mechanical properties, e.g. the low cycle fatigue strength.
A further problem of many known nickel-based superalloys, for example the alloys which are known from U.S. Pat. No. 5,435,861, is that in the case of large components, e.g. gas turbine blades or vanes with a length of more than 80 mm, the casting properties leave something to be desired.
The casting of a perfect, relatively large directionally solidified single-crystal component from a nickel-based superalloy is extremely difficult. Most of these components have defects, e.g., small-angle grain boundaries, freckles (i.e., defects caused by a series of identically directed grains with a high eutectic content), equiaxed limits of variation, microporosities, etc. These defects weaken the components at high temperatures, and consequently the desired service life and operating temperature of the turbine are not achieved.
However, since a perfectly cast single-crystal component is extremely expensive, the industry tends to permit as many defects as possible without the service life or operating temperature being adversely affected.
Another possibility is proposed in U.S. Pat. No. 7,115,175 B2: after the single-crystal component has been cast, the microporosities present, which were produced during the casting, are closed and eutectic γ/γ′ phase islands in the matrix are partially dissolved, in that a HIP (hot isostatic pressing) process is employed, then solution annealing for completely dissolving the eutectic γ/γ′ phase and for precipitating uniformly distributed large γ′ particles (referred to as octet shaped) is performed, and then precipitation heat treatment is carried out in order to obtain second and uniformly distributed fine cuboidal γ′ particles. This is intended to increase the strength of the superalloy.
According to the process described in U.S. Pat. No. 7,115,175 B2, the HIP process, which directly follows the casting step, is carried out after slow, two-stage heating of the cast object at a final HIP temperature in the range of 1174° C. (2145° F.) to 1440° C. (2625° F.), where the holding time is 3.5 to 4.5 hours and the pressure is in the range of 89.6 MPa (13 ksi) to 113 MPa (16.5 ksi), i.e., is relatively low.
This known process therefore produces single-crystal components made of nickel-based superalloys which are advantageously pore-free and have no eutectic γ/γ′ phases and have a γ′ morphology with a bimodal γ′ distribution.
It is not possible to positively influence the microstructure with respect to the undesirable rafting described above with the process disclosed in U.S. Pat. No. 7,115,175 B2.